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By G. T. Pruitt, T. L. Irby, P. B. Crawford, H. Ferrell

In water flooding peripheral, center-to-edge, line-drive or water-encroachment patterns the question has arisen, "when should a producing well be converted to a water-injection well?". It is realized that, if the producing well is shut in when water breakthrough occurs, some oil could be left in an unswept area. Yet, operation to high water-oil ratios may require processing much additional water. If a well is shut in, "where should the allowable be transferred?". A review of the literature indicates that very few data are available to the engineer to assist him in answering these questions. It is the purpose of this paper to present information which may be of value in selecting the flooding technique and the injection and producing wells and in scheduling injection and production rates. Certain phases of these problems can be studied in the laboratory. The Potentiometric model is well suited for studying many waterflooding problems. This model assumes that steady-state flow exists, that the mobility ratio is equal to one and that gravitational effects are neglected.'-' The three-phase sand model also has been recognized as an effective means for studying certain reservoir problems. Both models were used in making the study reported here. This paper discusses seven cases utilizing various waterflooding techniques. Data for the first five cases were obtained on the potentiometric model. Data for Cases VI and VII were obtained by constructing and flooding an unconsolidated sand pack (a three-fluid-phase system) and by determining the swept area and cumulative oil recovery for various types of operations. APPARATUS AND PROCEDURE The potentiometric model reservoir was 30 in. long X 10 in. wide by about 1/2 in. deep. Four wells were equally spaced in a line array along the lateral center line. (See Fig. 1.) Well No. 1 initially was an injection well for each of the patterns reported. The cases studied with the potentiometric model and sand model are described herein. The laboratory sand model reservoir was a 30- 10-in. rectangular Lucite vessel packed with sand to a depth of about 0.65 in. Four wells were equally spaced in a line array along the lateral center line. The reservoir material was an unconsolidated sand into which the reservoir fluids were mixed to yield fluid saturations of 55 per cent oil, 25 per cent water and 20 per cent gas. The water was added to the sand and the material mixed, after which the oil was added. Kerosene and water, the fluids used, gave a viscosity ratio of 1.64:1. In packing the reservoir, care was taken to assure constant fluid saturations and constant porosities. A new mixture was prepared for each run. During water flooding, a colored water was injected by a constant rate pump. The injection rate was 1.1 B/D/ft of sand thickness. For this series of tests, each of the producing wells was open to the atmosphere and, as in the potentiometric model study, Well No. 1 initially was an injection well for each of the patterns. KESULTS OF POTENTIOMETKIC MODEL STUDlES Two cases (Cases I and 11) involving equal pressure injections were studied, with the following results. Case Water injection was initiated at Well No. 1; the producing wells (Nos. 2, 3 and 4) were operated at constant and equal bottom-hole pressures. When water broke through at Well No. 2 and achieved the selected water-oil ratio, Well 2 was converted to a water injection well, resulting in two injection and two producing wells. The injection pressures at Wells 1 and 2 were maintained equal and constant; pressures at the producing wells were held equal, but they were lower than those at the injecting wells. A similar process was repeated when water broke through at Well No. 3. The flood was terminated when

By G. R. Purdy, J. A. Malcolm

The isothermal dendritic precipitation of y in p CuZnSn is compared with the isothermal formation of Widmanstatten a plates in p CuZn. Although both precipitates are evidently bounded by semicoherent interfaces, the dendritic shape of y evolves through smooth "sinuoidal" shape perturbations, while a-p interfaces tend to be strictly planar, or rnicroscopically faceted. An estimate of the average inter-facial free energy of cube-oriented y in p through the Mullins-Sekerka analysis indicates that it is low (of the order of 50 ergs per sq cm). Interface mication is apparently controlled by volunle dij'usion, and the observed anisotropy in growth probably results from anisotropy of inter facial free energy. Studies of the grain boundary precipitation of y in p indicate that the precipitate can grow in both $ grains, but becomes unstable with respect to the grain with which it shares a mobile semicoherent interface. Facets observed at the tips of Kurdjumov-Sacks-oriented Wid-manstatten a plates correspond to variants of an irrational hahit, near (231)0. Comparison with Hu and Smith's determination of relative interfacial free energies of a-(3 boundaries suggests that this system possesses an extremely deep o cup, which results in an ej'ective low-o zone. The possible effects of such a cup on the stability of a-p interfaces are discussed, a needles, grown during rapid continuous cooling, possess microscopically smooth interfaces, and are characterized by a high degree of' damage in the region of the boundaries. These boundaries, induced to move on tempering at 200"C, become faceted as equilibrium is approached, which seems to indicate that boundary structure and mobility in this system are dependent on the local degvee of departure from equilibrium MULLINS and sekerkal have elucidated the phase boundary stabilizing effect of surface tension. a factor not taken into account in previous treatments of morphological development. Their analysis applies to the isothermal growth of a pure spherical precipitate (planar in the limit of infinite radius) in a slightly supersaturated solution, under conditions of inter-facial local equilibrium and isotropy of interfacial free energy. More recently, several other factors which might tend to stabilize planar interfaces have received attention. These are: i) interface diffusion;4"8 and ii) slow interfacial reaction. cahn7 included a slight anisotropy in interfacial free energy, and showed that it has no stabilizing effect, but forces shape perturbations to reflect the anisotropy. shewmon4 noted that dendritic morphologies are rarely observed in solid-solid systems, and, in particular, postulated that slow interfacial reaction occurs at all boundaries between proeutectoid ferrite and austenite, except those which propagate parallel to {11 l)y planes (for Kurdjumov-Sachs [K-S] related phases). Previously, Aaronson~ had proposed interfacial reaction at the broad faces of Widmanstatten plates, and expected that the lateral motion of high-mobility incoherent ledge edges would produce planar precipitates bounded by law-mobility semi-coherent interfaces. The problem of interfacial reaction during the dif-fusional formation of precipitates has often been ignored, although it is apparent that a small component of chemical reaction must exist at any moving interface." Further, all solutions of the diffusion equations for boundary conditions corresponding to isothermal growth predict infinite growth rate at time zero, so that elements of interfacial reaction would be most easily detected at early times. However, reported deviations from diffusion control are comparatively rare," and the local-equilibrium diffusion-control model appears adequate for the description of many isothermal growth systems.'2"9 Cahn14 has treated the problem of migration of an interface which is able to reach positional energy minima, and concluded that such an interface, if subjected to a high local ther-modynamic "driving force", may migrate normal to itself without the benefit of ledges. As local equi-

Jan 1, 1968

By P. L. Raffo, T. E. Mitchell

Single crystals of Ta-Re alloys have been deformed in tension and compression as a function of solute concentration and temperature. The temperature dependence of the yield stress is decreased by alloying. At low temperatures there is a minimum in the yield stress (at about 2 at. pct) as a function of concentration. Three-stage hardening was not observed in the alloys. Linear hardening rates, where observed, were characteristic of stage II in pure tantalum. At low temperatures and high concentrations, work softening was often prominent at high strains. Electron microscopy revealed that dislocations are mostly of screw orientation, although heavily jogged and cusped, and often occurring in dipole pairs. The alloys become more prone to cleavage with increasing rhenium content, due mostly to an increasing yield stress rather than a change in cleavage strength. Alloying apparently decreases the Peierls stress of bcc metals at low temperatures by facilitating kink nucleation at solule atoms on screw dislocations. Higher solute concentrations further increase the yield stress. Quantitative calculations and comparisons with other bcc alloy systems are made to show that the predominant hardening mechanism is due to the interaction of dislocations with misfitting solute atoms close to their slip planes. The interaction is so localized that the core dilatation of dislocations could be important. CONSIDERABLE interest has arisen in recent years concerning the mechanical properties of unalloyed bcc metal single crystals.'-7 These and similar studies on polycrystalline materials have resulted in a considerable understanding of the macroscopic yielding characteristics8, 9 and the process of work hardening.'' This latter advance is due in a large part to the increased use of transmission electron microscopy in P I RAFFO junior Membar AlMF is Metallyureist Iewis such studies and the general observance of three stages of work hardening in the shear stress-shear strain curves. The properties of substitutional alloy single crystals of bcc metals have not been determined in as much detail as in the pure metals. Insufficient systematic studies have been made on the effect of alloying on the magnitude and temperature dependence of the yield stress or the stress-strain curves. The alloy systems which have been studied to some extent have been Ta-W,11, 12 Ta-Mo," Ta-Nb,12'14 Mo-Re,15 W-Re,16 and alloys of iron with nickel,17 platinum,17 vanadium,18 and aluminum.18 A large body of data exists on Fe-Si alloys, but these have mainly been a result of the experimental convenience of using this material in place of unalloyed iron, rather than a comprehensive study of the mechanical properties of the alloy system per se. The present authors have previously reported additional data on the properties of Ta-Re, Ta-W, and Ta-Mo alloy single crystals.1g Some of the data on Ta-Re alloys will be described later in this paper. Briefly, single crystals of Ta-Re (1.2 to 10.6 pct), Ta-W (6.2 and 10.2 pct), and Ta-Mo (2.9 and 9.0 pct) were grown by the floating zone method in vacuum. Compression tests were performed in the temperature range 77" to 623°K. Slip was found to occur in Ta-Re alloys at 77' and 298°K on the (~ll)[lil] system and slip lines were coarser and straighter than in unalloyed tantalum. No three-stage hardening was observed in any of the alloy systems and the curves were nearly parabolic. At temperatures above 175°K the yield stress* increased monotonically with in- creasing composition while at 77" and 175°K the yield stress initially passed through a minimum after which it began to increase again for concentrations greater than about 2 pct Re. Similar minima were observed in Ta-Mo alloys and to a lesser extent in Ta-W alloys. The minimum was interpreted as being a result of a

Jan 1, 1969

By Franklin J. Hill, Theodore D. Tiemann

Sized fractions of Wisconsin Gogebic taconite were reduced with hydrogen over the temperature range from 600° to 1000°C. In general, the degree and rate of reduction increase with temperature. Particle size has no observable effect except in the smaller fractions at 900°C and above, where reduction is impeded, possibly due to the siliceous nature of the ore. Gaseous diffusion appears to be significant as a rate controlling factor at certain stages. THE depletion of high grade iron ore reserves in the United States has established the need to utilize all available iron ore deposits, both in long range planning and from the viewpoint of national emergency. This has led to extensive research and development in this field. The beneficiation of Wisconsin Gogebic taconite ore has been previously studied by investigators at the University of Wiscconsin, 1 the United States Bureau of Mines,, Battelle Memorial Institute, 3 and others. Methods investigated have included concentration by flotation processes, magnetic and gravity separation, magnetic roasting and various combinations of them. The foregoing studies have been concerned with the production of an enriched ore suitable for blast furnace use. If the ore could first be reduced to metallic iron and then beneficiated, the product could be used directly in the open hearth or electric steel-making furnace. The by-passing of the blast furnace should enhance the economic feasibility of the beneficiating process. The work reported here,4 was undertaken to establish data on the hydrogen reduction of Wisconsin taconite that would eventually lead to the development of an economic direct reduction process for this and perhaps other low grade siliceous ores. The volume of work on direct reduction can be judged by the fact that 240 direct reduction processes were patented in the United States alone up to 1951, 5 and numerous additional processes have since been developed. Some of the factors contributing to direct reduction development are: 1) The increasing capital cost of the blast furnace and its accessory equipment. 2) The decreasing availability of high grade domestic ore. 3) The high cost of steel scrap for use in the open hearth and electric furnaces. In general, direct reduction processes can be divided into two types; those using a solid reductant (coal, peat, lignite) usually carried out in a kiln, and those using a gaseous reductant (H2, CO) in a shaft furnace or a fluidized bed retort. As there is no low cost coal readily available in northern Wisconsin but natural gas could be quite easily obtained, it is believed a gaseous reductant process would offer the greater feasibility for use on Wisconsin taconite. It is hoped that this preliminary work will beof interest to other investigators in this increasingly important field. TACONITE ORE The ore used in the investigation was obtained originally from the U. S. Bureau of Mines, and consisted of a representative composite made up from trench samples of the Norrie, Pabst, Plymouth, Pence, and Yale members of the Wisconsin Gogebic range iron formation. Chemically, the ore contains approximately 52 pct SiO, and 30 pct Fe. X-ray diffraction and other studies indicate the principal minerals to be quartz (SiO2), hematite (Fe2O3), and goethite (Fe2O3. H2O). Minor amounts of iron containing silicates, siderite (FecO3), and magnetite (FeO-Fe2O3) are present. The iron minerals and quartz are so finely disseminated that the ore requires grinding to less than 200 mesh (74 µ) for essential liberation.l,2 The iron content of each of the several size fractions studied is shown in Table I. EXPERIMENTAL METHOD The method used throughout the investigation was based on that developed by the U. S. Bureau of Mines in Minneaopolis6,7 in which the course of reduction is followed by the loss in weight of the ore sample. More specifically, a 100-g sample of the size frac-

Jan 1, 1962

By Clark L. Corey, Bogdan Lisowsky

Electrical resistivity, X-ray line positions, degree of order, and microstructures have been investigated for Ni-A1 alloys near the Ni3Al composition. The results indicate that Ni3Al undergoes disordering between 1250oC and the mp; also supersaturated y, produced in y + y' alloys by gas quenching, is extensively ordered prior to y' (Ni3Al) phase separation. The latter reaction, if not both, must follow the cooperative phenomena model. Shifts in line position as a function of quench-indulced super satulration are shown. Lattice pavarneters of equilibriunz and nonequilibrium phases are given. THE objective of this investigation was to determine the mechanism and kinetics of the reactions which occur in Ni3A1 and hypostoichiometric Ni3A1 alloys. The results are of particular significance relative to pre-precipitation phenomena. Early work on Ni-A1 binary compositions near the Ni3A1 composition (13.6 wt pct Al) has been adequately reviewed by Hanssen,1 whose phase diagram is presented in Fig. 1. williams2 verified the suggestion made by Taylor and Floyd3 and others4 that y', Ni3A1, may precipitate coherently on (100) planes of the y phase (see Fig. 1 for phase nomenclature); he further suggested a two-step precipitation mechanism consisting initially of probably short-range ordering followed by actual precipitation of y'. The precipitation stage in low-aluminum hypo-eutectic compositions has been studied by a number of investigators;*"7 short-range order or "zone" formation have been suggested as explanations for various diffuse X-ray scattering effects. The ordering reaction in Ni3A1 alloys is characterized by a) a very rapid rate of ordering in nonstoi-chiometric compositions, b) absence of a disordered phase region for the stoichiometric compound, y', or Ni3A1, c) precipitation phenomena as a dominant aspect in hypoeutectic alloys, and d) an increase in elec-trical resistivity on ordering.' The latter two factors were found by Newkirk to characterize ordering in Co-Pt, for which an extensive and detailed analysis was made.' An electrical-resistivity increase on ordering has also been reported for cuAu3.10 PROCEDURE Ni-A1 alloys in the form of 1300-g chill-cast ingots were prepared by vacuum melting electrolytic nickel and high-purity (99.99 pct) aluminum. Very small heats prepared from high-purity carbonyl nickel, supplied through the courtesy of Dr. Goff of the Naval Research Laboratory, were used for comparison with the larger ingots and no differences were noted. The cast material was homogenized in purified argon for a minimum of 72 hr at 1200°C and water-quenched from 1350°C. Compositions studied are given in Table I. The 10 pct A1 alloy was both hot- and cold-rolled. X-ray superlattice line intensities and peak positions were obtained from -270 mesh filings, using CuK, radiation with a LiF crystal monochromat adjusted to receive both 01 and a2. One-degree slits and a 1/8 deg per min scan rate were used to record line traces. A standard tungsten powder, obtained through the courtesy of G.E. Research Laboratory, was used to calibrate line positions.

Jan 1, 1968

By A. L. Lee, M. H. Gonzalez, R. F. Bukacek

Experimental viscosity data for methane are presented for temperatures from 100 to 340F and pressures from 200 to 8,000 psia. A summary is given of the available data for methane, and a comparison is presented for that data common to the experimental range reported in this paper. Correlation of the data is presented, and predicted values are given for temperatures up to 460F and pressures up to 10,000 psia. INTRODUCTION The increasing ranges of temperature and pressure at which fluids are produced, transferred and processed in the petroleum and chemical industries stress the need for accurate information on physical properties, both for engineering calculation and for improving the methods used to estimate physical properties. A survey of the literature reveals that disagreements between published data on the viscosity of methane are common and that most investigations have been conducted over restricted temperature and pressure ranges. This situation made it desirable to undertake an investigation with a scope common to most of the literature data available. APPARATUS AND MATERIALS The apparatus used in this investigation was the capillary tube viscometer'described by Dolan,10 with modifications introduced in the general design and operation of the instrument. It has an effective pressure range from 14.7 to 10,000 psia and a temperature range from room temperature to 400F. The design of the viscometer is based on the establishment of a manometric head between two vessels containing the test fluid and a volume of mercury. The reservoirs are connected by a capillary tube through which the test fluid flows and a tube through which mercury flows. A pressure gradient is established by elevating one of the vessels above the other, the resulting flow of mercury displacing the test fluid through the capillary. The schematic diagram of the system (Fig. 1) shows the arrangement of the equipment auxiliary to the viscometer. The density cells assembly (E, Fig. 1) was not used in this investigation since reliable data are available on methane. 23 The assembly consists of a bank of eight 316 stainless steel pycnometers. Experimental values of the density of a natural gas19 and i-butane14 samples have been obtained for pressures up to 8,000 psia and temperatures up to 340F. The methane used was obtained from the Southern California Gas Co. Mass spectrometric analysis showed a composition of 99.8 percent CH4, 0.1 percent C2H6 and 0.1 percent C3H8. EXPERIMENTAL VALUES AND ACCURACY OF DATA Determinations were carried out at several temperatures. With the temperature in the cabinet controlled at a given level, a series of runs were made at various pressures. To establish pressures lower than that obtained from the methane cylinder, it was necessary to vent the gas slowly to the atmosphere until the desired pressure was reached. The consistency of the test run time for a particular condition can be used as a criterion for the accuracy of the data. A further check is obtained by comparison of the viscosity values obtained using different driving heads. To establish the instrument's reproducibility, a large number of runs were made at three completely different conditions of temperature and pressure. Data obtained were analyzed statistically to determine the most reliable value for the variance of the experimental data. It was found that the variances of the data taken at different experimental conditions were not statistically equivalent, but they were small. The largest variance was found at 8,000 psia and 100F, where the standard deviation was 0.385 micropoise. Therefore, 99.9 percent of the data can be expected to fall within 1.15 micropoises of the population mean, except for the intervention of systematic errors in operation.

By A. E. Powers

YEARS of experience and research have shown that molybdenum, tungsten, and vanadium are among the most useful and effective elements in augmenting the high-temperature strength of heat-treatable, ferritic steels. Grün, for example, in conducting an early survey of the strengthening effects of various elements in annealed, low-carbon steel at 750° and 930°F (400° and 500°C), found molybdenum and vanadium to be the most effective.' Holtmann comprehensively studied the strengthening effects of molybdenum and vanadium in quenched and tempered steels at 930°F as a function of microstructure and carbon content.' He found creep strength to be related to the saturation of the austenite at the austenitizing temperature, such that strength is increased by alloying until the y loop is reached—the terminus of complete austenitization. Beyond this degree of alloying, the inability to heat treat will result in lower strength, at least for low and moderate creep temperatures. Tungsten, owing to higher cost than molybdenum, has been little used in ferritic high-temperature steels. As a result, very few investigations have been made of its high-temperature strengthening ability, and reliable data on its effectiveness are not available. Grün rated tungsten far less effective than molybdenum at 930°F (500°C) on a weight-percentage basis.' However, Tammann has found that tungsten is as effective as molybdenum in raising the recrystallization temperature of iron. For this reason Smith, in a review of the subject, has advocated a re-examination of the influence of tungsten on the high-temperature strength of steel.' In any study of the effects of alloying additions on high-temperature properties, recognition must be made of the many variables involved, some of which can be controlled but have been ignored in innumerable investigations of the past, and some of which are difficult or even impossible to control. One cannot, for example, assign a definite strengthening index to any one alloying element, for this will be dependent upon the microstructure (heat treatment and mechanical treatment), testing con- ditions (temperature, time, and stress), and the complete composition of the steel (accompanying alloying elements, impurities, etc.). One ever-appear ing variable, hardness or tensile strength at room temperature, may be eliminated by heat treating all of the alloys to a single hardness level. An attempt may be made to control the variable of primary quenched structure by oil or water quenching the test specimen. The control of such variables may not be possible in cases of widely varying compositions, but at least such deviations should be recognized. Material and Procedure The alloys, made as 30-lb ingots in an induction furnace, all had the same base composition of about 0.18 pct C, 0.85 pct Mn, and 0.48 pct Si. They were grouped as follows: Mo Group—Containing up to 5.2 pct Mo. W Group-Containing up to 6.0 pct W. V Group—Containing up to 3.3 pct V. Mo-V Group—Containing up to 2.7 pct Mo and up to 1.4 pct V. Mo-W Group—Containing up to 2.5 pct Mo and up to 2.5 pct W. The steels, listed in Table I, were forged to ¾ in. sq bars and the material for the high-temperature, rupture-test specimens was further swaged to %-in. round bars.

Jan 1, 1957

By Paul B. Crawford, Bobby L. Landrum

Electrical model studies have been made of the effect of drainhole drilling on production capacity. Drain holes are those lateral holes which are sometimes drilled to assist in increasing the oil drainage into the main wellbore. The studies indicate that if the drainhole length i.r near 0.2 of the drainage radius, the increase in production capacity may be approximately 1.8, 2.08, trnd 2.33 for the case of one, two and four drain holes, respectively. If multiple drain holes are drilled, shorter holes are required to achieve a specific increase in production capacity than if one or two holes are drilled. To achieve increases in production capacities by factors near three it is shown that the drainhole lengths divided by the drainage radii must be approximately 0.64, 0.48, cnd 0.37 for the case of one, two and four drain holes, respectively. IN T RO DUCTION Kecent developments in drilling techniques have permitted the drilling of lateral or horizontal holes to increase the rate of production. The number of these drain holes may vary from 1 to 10, and they have been drilled laterally as much as 100 ft. Drainhole drilling at substantial depths is of recent origin but has been used in the Spraberry field of West Texas, the La Pas field in Western Venezuela, and certain fields in California. The diameter of the drain holes has been approximately 6 in. In drilling drain holes it is desired to know how much increase in production capacity may be expected, what additional benefit may be gained by drilling more than one drain hole, and what increase in production capacity may be brought about by lengthening the existing drain holes. In realization of the importance of lateral drilling, electrical model studies were made to provide quantitative estimates on the effect of drain holes on production capacity. ANALYSIS An electrical model was used in studying this problem. A wooden circular reservoir approximately 30 in. in diameter and 2 in. deep was filled with a dilute salt solution to a depth of 1 in. A copper band formed the periphery of the tank, and a copper wire was inserted in the center of the tank to represent the well. The drainage radius divided by the well radius was equal to 500. A diagram of this type model is shown in the inset of Fig. 1. An electrical potential was applied to the wire well and copper band at the periphery to determine the flow of current. Following this measurement copper wires of various length were soldered to the vertical wire well in such a manner as to simulate lateral wells drilled into the reservoir. Multiple lateral holes of various lengths were studied by varying the number and length of copper wires soldered to the vertical wire well. All multiple lateral wells were located symmetrically about the vertical well. The effect on production capacity by drilling an infinite number of lateral holes was studied by replacing the lateral hole with a copper disk. For each condition the flow of current was measured and compared with the flow of the vertical well only. The model which is thus described provides the steady state increases in production capacity which may be expected in some cases. The data conform to those which would be observed for a uniform homogeneous reservoir. The data may have primary application to those reservoirs producing above the bubble point. It is realized that few lateral holes are drilled in reservoirs of this type, but the data may be of some value in guiding future work or aiding to explain observed phenomena.

Jan 1, 1956

INTRODUCTION Mineral explorationists range throughout the world looking for mineral deposits which can subsequently be made into mines. Mines are not found! Mines are made, usually at great cost and not often without significant risk. The process of making a mineral deposit into a mine involves a great deal more than a complete appraisal of the deposit. The making of a mine demands detailed feasibility studies of many factors, any one of which, if shown to be negative, could prevent the deposit from becoming a successful mine. Total feasibility analysis involves at least three major fields of study: (1) economic analysis, (2) financial analysis, and (3) political and environmental analysis. Economic analysis is essential for proper decision making at many stages in the evaluation of a mineral deposit. It is the yardstick by which the proposed mining project and alter- native investment opportunities are measured. Economic analysis is concerned with costs and potential profits and can be applied at various stages in the evaluation of a mineral deposit to determine if the project merits the additional funds necessary to bring about each succeeding stage. The nearer any deposit approaches to becoming a mine, the more carefully the economic analysis should be made. Financial analysis has reference to how, where, and at what cost the investment funds can be obtained to finance the proposed mine project. Political and environmental analysis has reference to the many intangible factors that can affect the making of a successful mine. These factors loom larger today than ever before, due to the profusion of government regulations and various environmental constraints now imposed upon the mineral industry both by law and by the weight of public opinion. Ultimately, many of these intangible factors can be resolved back into the economic analysis as an added cost to mine development, but more and more often they have become insurmountable obstacles to any development of certain deposits. The role of financial analysis and the analysis of intangible factors must never be underestimated when making a feasibility study upon which to base investment decisions; however, the emphasis of this text will be upon economic factors. THE PARAMETERS OF MINERAL ECONOMIC ANALYSIS The two major elements of economic analysis are: 1) The potential earnings or annual generated cash flow of the project. 2) The investment costs necessary to realize such earnings. If the difference between the two elements is large enough, i.e., if the return on investment from the cash flow is high enough to attract investors, the project can be said to be financially feasible. The mineral economic analysis must take into consideration many parameters, some of which are known or knowable within certain limits, and others of which are unknown or unpredictable within certain limits and for which estimates must be made. Because both known and estimated parameters enter into the same analytical computations, it is important for the analyst to realize the nature of each when considering the results of an analysis. Some parameters such as the grade of mining block, although estimated, are said to be known within mathematically defined limits. Other parameters, such as the future price for minerals produced, etc., can only be estimated or "guestimated" within reasonable limits for which no precision can be assigned. Known or Knowable Mineral Economic Parameters Known or knowable parameters are: 1) Tonnage and grade of valuable minerals for the total mineral body and for individual mining blocks. 2) Metallurgical recovery of valuable mineral. 3) Minimum operating life of mineral de- posit at any specified production rate. 4) Royalty schedule, property payments, etc. 5) Sales-purchasing agreements, smelter schedules, marketing contracts, etc.

Jan 1, 1980

By R. F. Domagala, M. Hansen, D. J. McPherson

On the basis of metallographic analysis, incipient melting data, thermal analysis work, and X-ray diffraction, phase relationships in the 0 to 50 atomic pct Cr region were carefully resolved. Phase relationships in the 50 to 100 atomic pct region were outlined. A single intermediate phase, ZrCr,, was established in the system. A eutectic and a eutectoid were found in the zirconium-rich region of the system, while a eutectic was located in the chromium-rich region. Limited solubility of chromium in both a and a zirconium was observed. THIS is another in a series of papers on zirconium-base binary phase diagrams based on a program sponsored by the Atomic Energy Commission. The Zr-Sn,' Zr-Si,' and Zr-Mo and Zr-W" systems have already been published. Previous work on the Zr-Cr system includes a partial diagram by Hayes, Roberson, and Davies' and a complete diagram by McQuillan." The phase relationships reported by these authors were determined with magnesium-reduced zirconium, whereas those reported herein are based on iodide metal. Some X-ray work on the intermediate phase in this system was also reported by Hayes et al., as well as by Wallbaum.' Comparisons with these data are made in the appropriate sections. Thermal analysis, metallographic and incipient melting techniques were employed to accurately resolve phase relationships in the 0 to 50 atomic pct Cr region, while the remainder of the diagram is outlined on a basis of cast structures and thermal analysis. A description and analyses of the metals used in the preparation of alloys for this study are included in Table I. A "low-hafnium" zirconium crystal bar (Westinghouse "Grade 3") was used. This material was produced by the decomposition of a volatile iodide onto a hot zirconium filament. The as-received crystal bar was coated with corrosion product from a standard autoclave test by which its grade designation is determined. A sand-blasting and subsequent HF-HNO, pickling treatment was employed to remove this contamination. The bars were cold-rolled to approximately 1/32-in. sheet, pickled again, sheared to 1/4-in. squares, washed with acetone, and stored for use. Ingot chromium metal from National Research Corp. and electrolytic chromium flakes from Johnson, Matthey and Co., Ltd., were available as alloy additions. The ingot chromium was received in granular form. The large granules were sized by crushing, treated for iron removal, rinsed with acetone, and stored. The flakes were broken to approximately 1/8 in. with mortar and pestle, cleaned with acetone, and stored for charging. Experimental Procedure A nonconsumable electrode arc furnace identical to that described by Hansen, Kessler, and McPher-son,' was used in this work. A 400 amp dc welding generator was the source of power. To insure homogeneity, ingots were melted in a water-cooled spun copper crucible, a total of four times without opening the furnace. No difficulties were encountered in the preparation of these binary alloys. Analyzed chromium contents were found to agree very well with the intended content. Melts of pure zirconium were interspersed in the course of preparing the alloys and their hardness was measured to keep a constant check on the melting technique. Vycor and quartz bulbs were used to contain the samples for isothermal annealing treatments. Vycor bulbs were employed for temperatures up to 110O°C, while quartz was used above this temperature. For treatments up to 950°C the bulbs were evacuated before sealing, but for higher temperatures a partial pressure of argon was admitted to the

Jan 1, 1954

By E. L. Kamen, H. D. Kessler, M. Hansen, D. J. McPherson

The highly reactive Ti-Mo and Ti-Cb alloys were prepared and heat treated under protective conditions. Phase diagrams were established based on micrographic and X-ray diffraction analysis and detection of incipient melting. ß titanium forms a continuous series of solid solutions with both molybdenum and columbium, whereas these metals are only slightly soluble in the a titanium modification. PREVIOUS work on the Ti-Mo and Ti-Cb systems was very limited and is discussed in succeeding sections of this paper relating to the respective phase diagrams. The diagrams in the present work were determined principally by micrographic analysis, because other methods such as thermal analysis and dilatometry are not suitable, due to the low diffusion rates of these alloys in the transformation range. X-ray analysis was believed less useful than micrographic analysis because: 1—On quenching certain high titanium alloys from above their transformation range, the body-centered cubic ß phase decomposes into acicular a (hexagonal close-packed) which room temperature X-ray diffraction techniques cannot differentiate from isothermal a. 2— High temperature X-ray techniques might be used, but are difficult to apply and are subject to error because of the high reactivity of titanium with gases such as oxygen and nitrogen. Although metallo-graphic study was the basic method used, X-ray diffraction analysis was applied to a limited extent, to further substantiate the findings. The transformation ranges on the titanium side of the Ti-Cb and Ti-Mo systems were studied both with DuPont Process A (magnesium-reduced) titanium alloys and with iodide titanium alloys. The Process A metal was used chiefly to bracket the transformation range in order to reduce the number of alloys given a particular treatment for the more accurate final study with iodide titanium alloys. Experimental Procedure The analyses of the metals used in the preparation of alloys for this study are presented in Table I. Melting Practice: Early in the work, a consumable electrode, arc melting furnace was used in the preparation of the Process A Ti-base Mo and Cb alloys. This type of furnace has been described by other investigators.', ' One of the major difficulties of the consumable electrode melting process was the prep- aration of homogeneous electrodes (with the mixture of coarse titanium sponge and fine powder alloying addition) which would give homogeneous ingots. Composition gradients encountered may have been due to a tendency for powdered alloying metals to drop away from the electrode in handling. This difficulty in the preparation of homogeneous alloys was overcome by the use of a nonconsumable electrode furnace for later work. This furnace is similar to the one used by Mc-Pherson and Fontana a Ohio State University. A diagram of the nonconsumable electrode arc furnace is given elsewhere.' The source of power was a 400 amp dc welding generator. Electrode tips of tungsten (% in. in diam) were used for all columbium melts and for a limited number of molybdenum melts. Molybdenum tips were used for the majority of molybdenum alloy melts. Columbium tipped electrodes were tried, but disintegrated rapidly at currents above 300 amp. In the Process A titanium melting practice the metals were charged into the water-cooled spun copper crucible in the following forms: l—titanium: sponge (—4 + 16 mesh); 2—molybdenum: sections of 0.003 in. sheet (approximately ½ in. sq); and 3— columbium: powder compacts (? in. lumps broken from large section). After alternately evacuating and purging the furnace several times with argon or helium, the arc was struck against the metal charge,

Jan 1, 1952

By Hsun Hu

Direct observations were made on the recrystalli-zntion of a cold-rolled (110) [001] crystal of 3 pct Si-Fe by annealing thin sections of the specimen inside the electron microscope during examination. Subgrain formation, presumably by polygonization, occurs rapidly, followed by a slow increase in the subgrain size before recrystallization can be detected. No migration of the subgrain boundaries was observed, and the formation of recrystallization nuclei by coalescence of subgrains is suggested. Re-c~ystallization characteristics of various texture components were individually examined. AS reported in a previous paper: orientations of the various deformation texture components in a 70 pct cold-rolled (110) [0011 crystal of 3 pct Si-Fe were completely confirmed by transmission electron diffraction patterns taken at various areas of the specimen. These orientations can be summarized as follows: Over most areas of the deformed crystal (111) planes are parallel to the surface. These areas correspond to either $- the two main defor: mation textures, the (111) [112] or the (111) [112].2-5 Within a few areas, (100) planes parallel to the surface are frequently found. These areas represent the minor deformation texture component, the (100) [Oll] type, which is derived from rotated Neumann bands.2'4'5 Only in very rare instances are (110) planes found parallel to the surface. This orientation corresponds to the very weak texture component of the (110) [001] type. In a pole figure, this (110) [0011 orientation is located within the orientation spread between the two main (111) [1121 type textures. Usually the electron diffraction patterns taken at various areas of the specimen are very simple ones, i.e., for any one area the pattern consists only of sharply defined spots of a single orientation. In some areas, however, more or less continuous rings of weaker intensities are superimposed on a simple pattern, indicating the presence of a wide range of additional orientations. In order to understand in detail the recrystallization of this cold-worked crystal, the recrywtallization characteristics of each of the various deformation texture components must be examined. This can best be done by using a high-temperature stage and transmission electron microscopy techniques, whereby small areas corresponding to the various texture components can be selected individually for examination. The present paper presents results obtained by this method. MATERIAL AND METHOD Specimens were prepared from a previously studied1 single crystal of a commercial grade 3 pct Si-Fe alloy with an initial orientation (110) [00l], cold-rolled 70 pct from 0.014 in. to 0.004 in. in thickness. Thin foils were prepared by electrolytic polishing as described in the earlier paper.' Specimens of this thin foil were then mounted in a high temperature stage similar to the one designed by Whelan.6 A direct current was used as the heating source. The temperature of the heating grid was calibrated by means of a Pt vs Pt-10 pct Rh thermocouple in a vacuum chamber (0.01 to 0.02µHg). Maximum variation within the temperature range of 250o to 000°C was about ± 25°C. After the desired areas were selected for examination, the specimen was heated to a low temperature for some time while the selected area was constantly watched. When no further change in structure and orientation could be observed, a higher temperature was applied. Photomicrographs and diffraction patterns were taken from time to time during examination, with the heater temporarily turned off* while the photographs were

Jan 1, 1962

By N. E. Garner

Laboratory experiments were conducted to investigate the cutting of a single diamond on limestone and shale under simulated down-hole conditions. A high-pressure chamber was modified by adding a rock-rotating device so that planing tests could be run. Weight on the diamond, diamond geometry and differential pressure across the face of the rock were varied. It was found that (I) high differential pressure reduced the volume of rock cut by the diamond at a given bit weight by strengthening the rock and changing the mode of failure; (2) a finely powdered rock left in the bottom of the grooves reduced the volume cut by as much as 50 percent; and (3) orientation of a single diamond about the axis of a drill point had considerable influence on diamond cutting efficiency. INTRODUCTION Diamond bits consist of many diamonds held in a matrix provided with water courses for fluid passage to clean, cool and lubricate the diamonds as they drill formation. Each diamond contributes its share to the overall effectiveness of a bit, but a lack of understanding of the performance of a single diamond has hampered efforts to engineer improvements in diamond bits. The need for those improvements prompted this study. This paper describes the findings of a laboratory investigation of the outting action of a single diamond where the differential pressure between the wdlbore and formation, diamond geometry and vertical force on the diamond were independently varied. The tests were conducted in a pressure chamber that provided relative horizontal motion (planing) between a rock and diamond under simulated down-hole pressure conditions. EQUIPMENT HIGH-PRESSURE CHAMBER WITH ROTATING DEVICE The high-pressure chamber used in previous rock mechanics studies' was modified for this study by addition of a rock-rotating device. The modified chamber (Fig. 1) differs from the old chamber in that the rock can be rotated to provide relative horizontal movement between the rock and the penetrator which may be a small element of a diamond or a drag bit. In this study the penetrator was a single diamond. The rock sample holder inside the chamber is mounted on a thrust bearing and is connected to the external drive mechanism by a stem that passes through two seal elements. The seal elements are needed to allow formation pressure to enter the bottom of the rock sample through a small hole in the stem. The Kapseal seal elements (Tefion boots backed by O-rings) provide low friction (hand free) even at differentials of 5,000 psi. The drive mechanism is a variable-speed transmission with a speed range from 0 to 30 rpm and an output torque of 2,800 in.-lb. A right-angle-drive gear box couples the drive to the stem. The new design has two independent pressure systems (formation and borehole) to simulate down-hole rock stress environment. Formation pressure enters from the bottom of the rock through a 6-in. diameter area, and borehole pressure acts over the remaining surface. A 14-in. layer of zero fluid loss oil-base mud poured over the top of the rock and an O-ring seal on bottom allow a differential between the borehole and formation pressures. Bayol 50 was used as the hydraulic fluid in the chamber. The differentials reported here were obtained by elevating the borehole pressure while keeping the formation pressure at zero.

By A. B. Parsons

NOT every organization on reaching the relatively ripe age of three score and fifteen can say with truth that its purpose and objects remain precisely the same as prescribed by its founding fathers. Of course, constancy of purpose may be something else than unalloyed virtue. It could merely reflect slavish worship of tradition, intellectual stagnation, or sheer inertia. On the other hand, it could establish the fact that the founders were men possessed of broad outlook, of prophetic vision, and of an acute sense of direction. In addition to charting the course with precision, the founders of the A.I.M.E. surely showed a flair for expressing clearly in a few words what others might say obscurely in many words. To quote verbatim from the "call" for the meeting, in 1871, the proposed organization was to have: "two great objectives: "First, the more economical production of the useful minerals and metals. "Second, the greater safety and welfare of those employed in these industries." This text underwent some refining on several subsequent occasions but today the ultimate product, as set forth in the Constitution reads: "to promote the arts and sciences connected with the economic production of the useful minerals and metals and the welfare of those employed in these industries by all lawful means." The original call was dated "Wilkes-Barre, Pa., April, 1871," and was signed by Eckley B. Coxe, of Jeddo, and by R. P. Rothwell and Martin Coryell, both of Wilkes-Barre, each of the three subscribing himself as "mining engineer." The authors proceeded to explain briefly just how it was proposed to achieve these objectives: "In European countries, where the arts of mining and metallurgy have long been the subject of the most careful study, no means have been found so effectual in attaining the end above proposed, as the free interchange of experience among those actually engaged in these industries; and this object has been accomplished mainly through the medium of `Institutes,' `Associations,' or `Societies,' composed of those engaged in these occupations, and by the periodical publication of `essays' or `papers' communicated to such societies by their members. "It must be evident to all practical men that the interchange of the varied experience

Jan 1, 1947

By H. M. Manasevit, R. L. Nolder, L. A. Moudy

The observation that silicon-on-sapphire crystallo-graphic relationships exist which are unrelated to one another implies that there are regions of sapphire orientations containing these relationships which are separated by boundaries. The presence of these boundaries is demonstrated, initially on rounded cylindrical surfaces md more definitively on flat sapphire substrates with orientations along primary crystallograPhic zones. By means of reflection electron and X-ray diffraction techniques the crystallographic relationships have been studied. Two new relationships have been found, viz., (111) Si (1014) sapphire and (001)Si (1012) sapphire. Also found were multicrystalline relationships. The approximated boundaries between the crystallographic regions for each relationship are shown by means of a stereo-graphic projection. These studies also clarify discrepancies in the literature concerning the silicon-sapphire relationships on the (1120) sapphire substrate. SINCE the first report of the deposition of single-crystal silicon on sapphire in 1963' this epitaxial system has been the object of study by many semiconductor laboratories. In particular, there has been strong interest in the crystallographic relationships between the two materials. In 19652,3 we demonstrated that at least four distinct modes of orientation relationships can occur, namely, (100)Si (0112) sapphire, (111)Si 11 (1124) sapphire, (111)Si (1120) sapphire, and (111)Si (0001) sapphire. Each relationship could be obtained when the substrate plane was at the particular orientation indicated. Other workers as well as our laboratory concurred in general with the initial results. However, two discrepancies were observed. First it was observed by Joyce et a1.4 that on a (1120) substrate a (110) silicon deposit could be obtained. This was actually consistent with the (111)Si (1124) sapphire mode. Second, two other laboratories,5,6 as well as our own,7 have found a deposition on the basal plane which is different than our initial findings. Even though a (111)Si deposit was obtained, the crystallographic relationships were not identical to those previously found. Subsequent work in our laboratories has confirmed that the form found in these later investigations of the basal plane is by far the more predominant, and the original finding was an anomaly. The relationships which have been previously estab1.ished are summarized in Table I. It is seen that at present a rather small number of possible substrate orientations has been investigated. Therefore, it seemed appropriate that a systematic investigation of silicon epitaxy on various sapphire substrate orientations would be of value. It was expected that the answers to three questions could be obtained from such an investigation. These are as follows: 1) Since we know of at least four distinct modes of orientation relationships, over what range of substrate orientation do these modes occur and what is the form of the boundary expected between the different regions in which each mode occurs? 2) Are there other modes which occur on substrate orientations not previously investigated? 3) What is the explanation for the discrepancy between our laboratory's findings and those results reported by Joyce et al.4 concerning the relationships occurring on (1120) substrates? APPROACHES AND EXPERIMENTAL METHODS Two different approaches were used. Initially, depositions were carried out on rounded cylindrical surfaces which lay perpendicular to the (0112) plane.* By this means the presence of boundaries and their approximate positions could be determined. Finally and more definitively, depositions on flat sapphire substrates with a variety of orientations, mostly along primary crystallographic zones and about the (1120) plane, were studied by means of reflection electron diffraction and X-ray diffraction (using back-reflection Laue patterns and full circle goniometer techniques) in order to determine crystallographic relationships. The sapphire orientations investigated by these means

Jan 1, 1969

By D. R. Mash

A strong yield point with attendant enhanced mechanical properties was found in commercially pure beryllium under certain conditions of heat treatment. Beryllium specimens also responded to both quench and strain aging treatments. Based on these results, a possible means of obtaining increased normal temperature mechanical properties is suggested. COMMERCIALLY pure beryllium has been found to exhibit the yield point phenomenon and several other aging manifestations under certain conditions of heat treatment. This behavior is not surprising in view of the present accumulation of experience regarding such effects in numerous metals; but in addition to extending this experience to include beryllium, the present study is of interest because aging effects are generally associated with the presence of small amounts of solute elements. The well known lack of tensile ductility in beryllium at normal temperatures is also considered to be due to impurities in some way. In the present investigation, the view was taken that beryllium of commercial purity (greater than 98 pct) should be considered as a dilute alloy and that precipitation phenomena may well exist which could lead to enhanced mechanical properties. In connection with this possibility, published data on the variation of mechanical properties with temperature are of interest. For example, Fig. 1 illustrates the variation of ductility (percentage of elongation) and ultimate strength as a function of temperature for beryllium of various histories.* Although these curves are somewhat schematic in nature, they clearly indicate the possibility that precipitation of an unstable phase and/or aging effects occur during tests in the region 400" to 700°C. The existence of such behavior should be recognized as affecting the properties of the commercial material and perhaps can be employed advantageously. Experimental Method All testing was done at room temperature using a Baldwin universal testing machine. A microformer extensometer was used in conjunction with a micro-former-type stress-strain recorder. All tests were conducted at constant loading rate, the slowest rates being used. This corresponded to 120 lb per min in the 1200 lb full scale range and 300 lb per min in the 6000 lb full scale range. After the rate of loading was set in the elastic range of the material during the early part of the test, no further adjustments were made, allowing the test to proceed at that setting. In the plastic region of the curve above the yield region, strain rates of from 0.002 to 0.006 in. per min were regularly obtained. All experimentation was carried out on specimens machined from the same lot of beryllium, lot No. Y5036BP. Table I shows the chemical analysis as furnished by the supplier, Brush Beryllium Co., as well as the standard specifications to which commercial beryllium is produced. Both 1/4 in. thick hot pressed plate and ½ in. round hot pressed and warm extruded rods were obtained. The history of the material is reported to be as follows: 1—Beryllium metal plates, produced from a standard hot pressed QMV beryllium metal block 4x15x40 in., were cut off and finished machined to 1/4x6x12 in. flat plates. The original hot pressed block was made by hot pressing —200 mesh vacuum cast attri-tioned powder at about 1050°C under vacuum. 2—Beryllium rod was hot pressed and warm extruded. Extrusion billets were machined from the regular hot pressed blocks followed by extrusion at 425°C. The extrusion was done without jacketing, and graphite was used as a lubricant. The rods were then straightened at 720 °C, after which they were annealed at 800°C for 15 min. Both flat tensile specimens and round tensile specimens were tested. The former were machined from the plate stock into standard specimens according to specifications shown in ASTM E8-52T. The nominal cross-section was 0.25x0.500 in. and the gage length was 2 in. The round specimens were

Jan 1, 1956

By F. Weinberg

In developing a mechanism for the solidification of metals from the melt, it has been proposed that solidification proceeds by the growth of platelets parallel to close packed planes. The evidence for this is based primarily on observations of decanted interfaces of lead1,2 and tin on which platelets 1 to 10 µ thick4 were observed. In recent observations of decanted eutectic alloys, 5,6 it was found that the eutectic structure on the decanted surface was markedly different than that of the bulk material, with the surface structure extending approximately 10 µ behind the interface.' This suggested that, at least for these experiments, a residual liquid layer 10 µ thick had been left on the decanted interface, which solidified subsequent to decanting. If a similar thickness of residual liquid was left on the decanted surfaces in the lead and tin observations referred to above, then it is questionable to what extent the observed platelets can be related to the solid-liquid interface before decanting. The purpose of the present investigation was to measure the thickness of the residual liquid layer on a decanted surface of tin, using relatively standard decanting procedures. The technique used was to add approximately 200 ppm of Tl204 (PI) to the melted portion of a high purity (99.999 pet) bar of tin, which was being progressively melted from one end, and then decant during melting. Since only the melt contained the radioactive tracer, the activity emanating from the decanted surface could be used as a measure of the thickness of the residual liquid layer. Progressive melting was carried out in the same fashion as single crystal growth, namely, in a horizontal graphite boat, in an argon atmosphere, using a moving furnace. The decanting was done in two ways; 1) removing the furnace and suddenly dropping the melt end of the boat, and 2) suddenly jerking the solid away from the liquid by the action of a falling weight (estimated peak acceleration 3 x 105 cm per sq sec), using a split graphite container. In all cases the melt was vigorously agitated by bubbling argon through it. The activity of the decanted surface was measured through an aperture 1.5 cm by 0.5 cm in a lead shield with an Anton tube geiger counter. Assuming that the active layer on the surface had the same concentration of Tl204 as the decanted liquid, the thickness of the layer could be determined from the measured activity of the decanted material, and the relationship between activity and material thickness. The latter was determined by rolling some of the decanted material into foil and measuring the activity of multiple layers of the foil. The activity At of material of thickness t (in mm) was determined to be A, = Ao(1 - e-22.5t) where Ao is the activity of the bulk material. The results of the measurements of the thickness of the residual liquid layer for six decants are given in Table I. The variation in t is attributed to variations in the decanting procedure and the rate of melting, the latter being difficult to control because of the vigorous agitation of the liquid. The rate of melting was in the order of 2 cm per hr, estimated from the position of the solid-liquid interface during melting. The results indicate that a residual liquid layer of the order of 10 µ in thickness is left on the decanted interface in agreement with the observation on eutectic alloys referred to above. If the mixing of the liquid during melting is not complete, as was

Jan 1, 1962

By E. S. Bumps

A series of macrographs are presented for the purpose of graphically recording the strain phenomenon that accompanies the energy transition in the Charpy impact testing of a mild steel that is susceptible to strain aging. STRAIN patterns in mild steel Charpy impact specimens were obtained in order to determine whether, on the first appearance of plastic strain during deformation, the initial fracture cracks appear in the region of plastic strain or in the elasti-cally deformed regions between. Although this attempt to locate the initiation of cleavage was not successful, the observed strain patterns illustrating various stress conditions in partially deformed Charpy specimens are presented along with a discussion of the possible dependence of transition energy on the yield phenomenon. Earlier work on the development of strain patterns in mild steel have been reported by Fry1 and later by Jevon.2 Harris, Rinebolt, and Raring3 observed characteristic strain markings on the surface of Charpy V-notch impact specimens fractured above and below the transition temperature which they presented schematically. It is intended that the strain figures presented here will add to the literature a graphic correlation of strain with transition energy in impact specimens. Samples of % in. round hot rolled, 0.20 pct C steel were machined to standard 0.394 in. sq, V-notch Charpy impact specimens. Preliminary testing involved determination of transition temperature, Fig. 1, using a Sontagg impact tester. These data were subsequently used to select temperatures above, within, and below the transition range at which specimens were partially deformed. Partial deformation was obtained by either varying the energy or the testing temperature, see Table I. Macroscopic strain figures shown in Figs. 2 to 19 were produced by several metallographic repolish-ing and etching steps. The macroetchant used was as follows:4 120 g HCl, concentrated; 100 ml water; and 90 g cupric chloride. Strain figures were obtained on the outer surface and on sections through the central portion of the impact specimens in order to determine an approximation of the volumetric strain pattern. In the tensile strain figures reported by Jevon, it was found that strain etching would not work immediately following strain unless the material was heated to 200°C. The impact strain patterns presented here were obtained without this heating, but the specimens were not etched until ten months after straining. This confirms a discussion of Jevon's paper, in which it was pointed out that strain figures can also be developed after aging for long periods of time at room temperature. Hardness of the strained and unstrained portions, revealed by the etched pattern, was explored using R, impressions converted to R, hardness values. Discussion of Results In the impact test, as in the tensile test, some parts of the specimen yield and deform plastically, while other regions are deformed only elastically. The plastic deformation bears a relation to the imposed stress pattern. In the strain figures presented here, the plastically deformed regions are the dark etching ones. The strain figures are always symmetrical because the stress state in the Charpy-type specimen under impact loading is symmetrical. The difficulties encountered to produce initial cracks are intimated in Figs. 2 and 4. At a temperature of —75ºF, the stress was not sufficient to pro-

Jan 1, 1953

By A. B. Cummins

THE past year has been of unusual interest for industrial minerals. It is increasingly evident that requirements for these raw materials move with general economic trends. Thus, with a peak year in the construction, chemical and petroleum industries, the nonmetallic mineral industries have been brought into the complex maze of an economy in which military spending is in the billions and in which the national income and wholesale commodity prices reached new levels. Many industrial minerals are of first rank essentiality in the defense effort and its supporting industries, and since some of these are of limited supply and others dependent on imports, the problems of procurement, distribution and conservation of available supplies make a complicated pattern. Despite some drawbacks and weak spots, the industrial mineral industries have, in general, met the national situation with great credit and many new production records have been set. The general accomplishment in 1951 should be one of satisfaction with confidence for the future. Difficulties, unsolved problems and uncertainties should not obscure the general trend which many consider optimistic. Prices in the main have been upward, in line with the inflation spiral, but nonmetallics appear with some exceptions to be behind the general increases. Labor difficulties have not been serious throughout the year. Retarding influences in expansions and planning have been the uncertainty of the Korean and international situations and a lack of confidence by many in the government economic program and its ultimate impact on the economic structure. It must also be kept in mind that stockpiling, unless all-out war develops, reaches a level and then ceases to be a major sustaining factor in the mineral markets. In accordance with the Defense Production Act of 1950 various control agencies have been set up to establish the desired rate of stockpile acquisitions. The objectives are subject to change depending on availabilities and needs-military, industrial, and civilian. In some minerals there has been disagreement on actual stockpile requirements, and there is some feeling that the needs of defense might be better met if more funds went into the development of new sources of supply of scarce raw materials and more materials were released to industry at this time. Nonmetallic minerals on the strategic and critical list of the Munitions Board as of July 1951 were: Group I; asbestos, some types; bauxite, refractory grade; beryl; celestite; chromite, chemical and refractory grades; columbite; corundum; diamonds, industrial; fluorspar, acid grade; graphite, some types; kyanite; manganese ore, battery and chemical grades; mica, some types; quartz crystals; rare earths; sapphire and ruby; talc, steatite, block; and Group II; bauxite, abrasive; cryolite, natural; diamond dies; graphite, crystalline fines; iodine; mica, muscovite block, stained and lower; rutile; selenium; talc, steatite, ground; zirconium ore, zircon. Details on the current status of stockpiling inventories are not available. The Defense Minerals Exploration Administration has, since April, executed 39 contracts for the exploration of nonmetallic minerals, including twenty for mica, seven for beryl-mica, six for asbestos, three for sulphur and one each for monazite, steatite talc and fluorspar. Total value of these contracts came to $892,059, of which the Government's participation amounts to $689,247. There were on hand at year's end another 97 applications for exploration aid on nonmetallic minerals in twenty different

Jan 1, 1952

By H. R. Ogden, E. S. Bartlett, A. G. Imgram

Notched and unnotched tensile specimens of wrought and recrystallized, oxygmted and hydro-genated tantalum and columbium were tested over a range of temperatures selected to encompass the ductile-to-brittle transition. Increasing the oxygen content of tantalum and columbium increased their strength, and raised their ductility transition temperature, but did not cause a marked deterioration of notch properties below that of commercially pure material. The adverse effects of oxygen were more pronounced for recrystallized structures than for wrought material. The effects of hydrogen are complicated by strain aging or slow strain rate hydrogen embrittlement at -105°F. However, in general, hydrogen had only a minor effect an strength, but markedly increased the effective transition tem-terature and caused a deterioration of notch properties . BCC metals including the refractory metals in Groups Va and VIa (V, Cb, Ta and Cr, Mo, W) can be expected to undergo a ductile-to-brittle transition and exhibit brittle fractures at low temperatures.' Notch embrittlement is associated with this phenomenon, and can be a serious design problem. Table I compares the tensile transition behavior and notch sensitivity of commercially pure Group Va and VIa refractory metals. The Group VIa refractory metals have considerably higher transition temperatures, and become notch sensitive at higher temperatures, than do the Group Va refractory metals. In addition, the Group VIa refractory metals are sensitive to metallurgical structure. Recrystallized structures generally have higher transition temperatures than wrought material of the same composition. The relatively poor transition behavior of the recrystallized structures of these metals is generally attributed to the segregation of interstitial impurity atoms at grain boundaries. Pronounced sensitivity to structure is not evident with commercially pure Group Va refractory metals, however. Since these metals have higher solubilities for interstitial elements than do the Group VIa metals ,4 it can be expected that higher total interstitial concentrations are necessary to develop a comparable degree of grain boundary segregation and to produce a similar sensitivity of transition temperature to structure. Increasing only the concentration of interstitial atoms in solution would tend to increase the flow stress of a metal in relation to its cleavage fracture stress and thereby raise its transition temperature, but would not be expected to cause a marked sensitivity of transition temperature to structure. A research program has been conducted at the Battelle Memorial Institute to evaluate the effects of interstitial oxygen and hydrogen on the transition behavior and notch sensitivity of tantalum and columbium, and to evaluate the effects of structure on "contaminated" material. EXPERIMENTAL PROCEDURES Notched and unnotched tensile specimens of oxygenated and hydrogenated tantalum and columbium with both wrought and recrystallized structures were tested over a range of temperatures selected to encompass the ductile-to-brittle transition Oxygen levels selected for study in both tantalum and columbium were nominally 500 and 1000 ppm, whereas nominal hydrogen contents were 150 and 500 ppm in tantalum and 150 and 300 ppm in columbium. Oxygen and hydrogen additions were made directly to 6-in. lengths of 1/2-in.-diam electron beam melted wrought bar stock, in a Sieverts Apparatus. Subsequent high-temperature annealing was required for homogenization of the oxygenated material, but was not necessary in the case of hydrogenated stock. The oxygenated and hydrogenated bars were then swaged to 1/4-in.-diam rods. Tensile specimen blanks were cut from the 1/4-in.-diam rods and annealed to achieve stress-re-

Jan 1, 1963